Ni-Based Superalloy Powder for Additive Manufacturing and an Article Made Therefrom

ABSTRACT

A nickel base superalloy powder for additive manufacturing applications is disclosed. The alloy powder has the following broad weight percent composition:C 0-0.1Mn0.5 max.Si  0-0.03Cr 4-16Fe 0-1.5Mo0-6W0-8Co 0-15Ti0-2Al0.5-5.5Nb0-6Ta 7.5-14.5Hf 0-2.0Zr 0-0.1Re0-6Ru0-3B  0-0.03The balance of the alloy is at least 50% nickel and the usual impurities. An article of manufacture made from the alloy is also disclosed.

CROSS REFERENCE TO RELATED APPLICATION

This application claims the benefit of U.S. Provisional PatentApplication No. 62/795,618, filed Jan. 23, 2019, the entirety of whichis incorporated herein by reference.

BACKGROUND OF THE INVENTION Field of the Invention

This invention relates generally to Ni-based superalloys and inparticular to a Ni-based superalloy powder that is useful in an additivemanufacturing process for making articles therefrom.

Description of the Related Art

With the advent of additive manufacturing technologies, manufacturers ofturbomachinery (high temperature gas turbines and jet engines) arelooking at leveraging the design flexibility offered by the additivemanufacturing process to directly produce parts with complex featuresfor use in the hot gas path of gas turbines and jet engines. A powderbed fusion technique using selective laser melting is one of the mostpromising additive manufacturing techniques. Compared with the knownelectron beam powder bed process, selective laser melting with powderbed does not require pre-sintering. Consequently, designers canincorporate complex internal features in the parts without worryingabout removing sintered powder trapped inside hollows of the part afterprinting. Laser systems also offer finer resolution for printingintricate features.

The drawback of laser melting system is the large temperature gradientthat occurs in a part during solidification. The very large temperaturegradient leads to cracking of the alloy, especially for superalloyshaving a high volume fraction of the gamma prime (γ′) phase thatprovides superior strength and creep resistance such as CM247LC andRene142. For selective laser melting process, cracking occurs eitherduring build or post-build heat treatment. Efforts have been made by thesuperalloy research community to mitigate the cracking problem, but havefocused primarily on adjusting laser parameters such as power andscanning speed. There have been few attempts to modify the chemistriesof the known superalloys, such as described in US20180347014A1, becauseof the reluctance to deviate from the known and accepted alloy chemistryspecifications.

SUMMARY OF THE INVENTION

The problems associated with the known superalloys as discussed aboveare resolved to a large degree by an alloy powder having the followingbroad and intermediate weight percent ranges.

Element Broad Intermediate C  0-0.1  0-0.1 Mn 0.5 max. 0.5 max. Si  0-0.03   0-0.03 Cr  4-16  4-11 Fe  0-1.5 0-1 Mo 0-6 0-5 W 0-8 0.5-8 Co  0-15  0-10 Ti 0-2 0-1 Al 0.5-5.5 3-5 Nb 0-6 0-5 Ta  7.5-14.5 7.5-11.5 Hf  0-2.0  0-1.5 Zr  0-0.1  0-0.1 Re 0-6 0-5 Ru 0-3 0-3 B  0-0.03   0-0.03

The balance of the alloy is at least about 50% nickel and the usualimpurities found in commercial grades of Ni-base superalloys intendedfor the same use and service. The alloy is further characterized by thefollowing relationships among the constituent elements of the alloy:

-   -   % Al+% Ti≤6;    -   % W+% Mo+% Ru+% Re<10; and    -   % C+% B+% Si+% Zr<0.15.

In accordance with another aspect of the present invention, the Ni-basesuperalloy powder may be embodied with any of the following preferredalloy weight percent compositions.

Pre- Pre- Pre- Pre- Pre- ferred ferred ferred ferred ferred Element A BC D E C  0-0.1  0-0.1  0-0.1  0-0.1 0-0.1 Mn 0.5 max.   0.5 max.   0.5max.   0.5 max.   0.5 max.   Si   0-0.03   0-0.03   0-0.03   0-0.03 0-0.03 Cr  9-11  9-11 4-8  9-11 8-10  Fe 1 max. 1 max. 1 max. 1 max. 1max Mo 2-5 2-5 0.5-1.5 2-5 2-5  W 0.5-2  0.5-2  6-8 3.5-4.5 4-6  Co 4-60.5-2  0-1 0.5-2  9-11  Ti  0-0.5  0-0.5 0-1  0-0.5 0-0.5 Al 3-5 3-5 3-53-5 4-5.5 Nb 3-5 2.5-4  0-1 2.5-4  0-1  Ta 7.5-9.5  8.5-10.5  8.5-10.5 8.5-10.5 9.5-11.5  Hf  0-0.5  0-0.5 0.5-1.5  0-0.5 0-0.5 Zr  0-0.1 0-0.1  0-0.1  0-0.1 0-0.1 Re 0-1 0-1 0-1 0-1 0-1  Ru 0-1 0-1 0-1 0-10-1  B   0-0.03   0-0.03   0-0.03   0-0.03  0-0.03

The balance of the alloy is at least about 50% nickel and the usualimpurities found in commercial grades of Ni-base superalloys intendedfor the same use and service. The alloy is further characterized by thefollowing relationships among the constituent elements of the alloy:

-   -   % Al+% Ti≤6;    -   % W+% Mo+% Ru+% Re<10; and    -   % C+% B+% Si+% Zr<0.1.

The foregoing tabulations are provided as a convenient summary and arenot intended to restrict the lower and upper values of the ranges of theindividual elements for use in combination with each other, or torestrict the ranges of the elements for use solely in combination witheach other. Thus, one or more of the ranges can be used with one or moreof the other ranges for the remaining elements. In addition, a minimumor maximum for an element of one range can be used with the minimum ormaximum for the same element in another range, and vice versa. Moreover,the alloy according to the present invention may comprise, consistessentially of, or consist of the constituent elements described aboveand throughout this application.

The alloy according to the present invention is designed to address thefollowing problems: the transformation strain associated with theprecipitation of γ′ during cooling and the internal stress buildup thatoccurs during aging heat treatment. Those phenomena combined with thefast γ′ precipitation kinetics and the residual stress inherent in theadditive manufacturing process make the known alloys prone to strain-agecracking defects. It is an objective of the present invention to modifythe known alloy chemistry and thereby, to tune the lattice misfit andprecipitation kinetics to provide a new alloy that is optimized for theselective laser additive manufacturing process and which is strain-agecracking resistant while retaining a high γ′ volume fraction forsuperior elevated temperature strength and creep resistance.

In accordance with another aspect of the present invention there isprovided an article made from consolidated powder produced with any ofthe alloys described above. The article is characterized by having morethan about 35 volume percent (vol. %) of the γ′ phase in gamma (γ) phasematrix when the article is in the age hardened condition. The article isfurther characterized by a lattice misfit parameter (δ) that is greaterthan about −0.1% and preferably about −0.05% to about +0.6% where thelattice misfit parameter is defined as:

$\delta = \frac{2 \times \left( {\alpha_{\gamma^{\prime}} - \alpha_{\gamma}} \right)}{\alpha_{\gamma^{\prime}} + \alpha_{\gamma}}$

The parameter α_(γ′) is the lattice constant for the γ′ precipitate andparameter α_(γ) is the lattice constant for the γ matrix material.

Here and throughout this specification the term “percent” or the symbol“%” means percent by weight or mass percent, unless otherwise specified.The term “high γ′ volume fraction” means at least 35% by volume (volume% or vol. %) of the γ′ phase in gamma (γ) phase matrix of the alloy whenin the age hardened condition. The term “high strength” means a yieldstrength higher than that provided by Alloy 718 when produced with anadditive manufacturing process tested at a temperature beyond 1500° F.(816° C.). The term “high creep resistance” means a creep strengthhigher than that provided by Alloy 718 produced with an additivemanufacturing process tested at a temperature beyond 1500° F. (816° C.).The term “aging” is used synonymously with “age hardening” and“precipitation hardening”. Moreover, the term “solvus” means the solvustemperature as that term is generally understood.

BRIEF DESCRIPTION OF THE DRAWINGS

The working examples described in the following detailed description ofthe invention will be better understood when read with reference to thedrawings wherein:

FIG. 1A is a backscatter SEM image of an aged sample of Example 1 of thealloy according to the present invention;

FIG. 1B is a backscatter SEM image of an aged sample of Example 2 of thealloy according to the present invention;

FIG. 1C is a backscatter SEM image of an aged sample of Example 3 of thealloy according to the present invention;

FIG. 1D is a backscatter SEM image of an aged sample of the CM247LCalloy;

FIG. 2 shows a graph of the calculated misfit parameter compared toactual measured values of the misfit parameter for Examples 1-3 and theCM247L alloy;

FIG. 3A is an optical micrographic image of a sample of Example 1 afterexposure at 1600° F. (871° C.) for 1000 hours showing an oxide layer atthe surface of the sample;

FIG. 3B is an optical micrographic image of a sample of Example 2 afterexposure at 1600° F. (871° C.) for 1000 hours showing an oxide layer atthe surface of the sample;

FIG. 3C is an optical micrographic image of a sample of Example 3 afterexposure at 1600° F. (871° C.) for 1000 hours showing an oxide layer atthe surface of the sample;

FIG. 4A is an optical micrographic image of a sample of Example 1 of thealloy according to the present invention showing a resolidified meltpool after laser melting of the metal;

FIG. 4B is an optical micrographic image of a sample of Example 2 of thealloy according to the present invention showing a resolidified meltpool after laser melting of the metal;

FIG. 4C is an optical micrographic image of a sample of Example 3 of thealloy according to the present invention showing a resolidified meltpool after laser melting of the metal;

FIG. 4D is an optical micrographic image of a sample of the CM247LCalloy showing a resolidified melt pool after laser melting of the metal;

FIG. 5A is an optical micrographic image of a sample of Example 1 of thealloy according to this invention after laser melting and an aging heattreatment;

FIG. 5B is an optical micrographic image of a sample of Example 2 of thealloy according to this invention after laser melting and an aging heattreatment;

FIG. 5C is an optical micrographic image of a sample of Example 3 of thealloy according to this invention after laser melting and an aging heattreatment;

FIG. 5D is an optical micrographic image of a sample of the CM247LCalloy after laser melting and an aging heat treatment;

FIG. 6A is a series of optical micrographic images of samples frombuilds of CM247LC alloy powder prepared with a laser powder bed fusionsystem at various laser power densities and various hatch spacings;

FIG. 6B is a series of optical micrographic images of samples frombuilds of alloy powder made from Example 1 and prepared with the laserpowder bed fusion system at various laser power densities and varioushatch spacings;

FIG. 6C is a series of optical micrographic images of samples frombuilds of alloy powder made from Example 2 and prepared with the laserpowder bed fusion system at various laser power densities and varioushatch spacings;

FIG. 7A is an optical micrographic image of one sample of the as-builtCM247LC alloy after solution and aging heat treatments;

FIG. 7B is an optical micrographic image of one sample of the as-builtalloy of Example 1 after solution and aging heat treatments;

FIG. 7C is an optical micrographic image of one sample of the as-builtalloy of Example 2 after solution and aging heat treatments;

FIG. 8 shows graphs of the tensile and yield strengths of Example 4 andExample 5 as a function of temperature in the range from roomtemperature to 1800° F. (982° C.);

FIG. 9 shows graphs representing the stress rupture lives of Example 4and Example 5 based on the Larson-Miller parameter; and

FIG. 10 shows graphs of the results of the 1800° F. (982° C.) cyclicoxidation testing of Example 4 and Example 5.

DETAILED DESCRIPTION

The alloy of this invention contains at least about 0.5% and preferablyat least about 3% aluminum to combine with available nickel to form theγ′ phase, Ni₃(Al,Ti), during cooling after laser melting and duringsubsequent aging of the alloy in consolidated form. The alloy containsnot more than about 5.5% and preferably not more than about 5% aluminumin order to benefit the weldability of the alloy. Up to about 2%titanium can be substituted for some of the aluminum for similarreasons. Preferably, the alloy contains not more than about 1% and forsome applications not more than about 0.5% titanium. The combined amountof Ti and Al should not be more than about 6%. Preferably, the alloy maycontain about 3% to 6% and better yet, about 3.5% to 5% of aluminum andtitanium combined (Al+Ti).

This alloy also contains at least about 7.5% tantalum. However, in apreferred embodiment the alloy may contain at least about 8.5% tantalum.Tantalum also combines with nickel to form the γ′ phase during coolingafter laser melting and during subsequent aging of the alloy inconsolidated form. Too much tantalum can lead to the formation ofundesirable phases such as eta phase and delta phase. Therefore,tantalum is restricted to not more than about 10.5% and preferably tonot more than about 9.5%. In a different embodiment, the alloy containsabout 9.5% to about 11.5% tantalum. Up to about 6% niobium and up toabout 2.0% hafnium may be present in this alloy in substitution for someof the tantalum for the same purpose. Tantalum, and niobium and hafniumwhen present, ensure that a sufficient volume of the γ′ phase is formedduring the aging heat treatment of the additively manufactured product.

The alloy may optionally contain up to about 6% molybdenum, up to about8% tungsten, up to about 6% rhenium, and up to about 3% ruthenium. Thoseelements tend to partition to the matrix material and theirconcentrations are controlled to shift the lattice misfit parametertoward a more positive value. A more positive lattice misfit parameteris believed to mitigate strain-age cracking that could otherwise occurin the alloy during solidification and subsequent aging heat treatment.When present, molybdenum is preferably limited to not more than about5%, rhenium is limited to not more than about 5%, ruthenium is limitedto not more than about 1%. For certain applications molybdenum islimited to not more than about 1.5%, tungsten is limited to not morethan about 2%, and rhenium is limited to not more than about 1%. Whenpresent in this alloy, the combined amounts of molybdenum, tungsten,rhenium, and ruthenium are controlled such that % Mo+% W+% Re+% Rh isless than 12%, preferably about 2% to about 12%, and better yet about 3%to about 10%.

This alloy may contain up to about 15% cobalt or up to about 11% cobalt.Cobalt lowers the stacking fault energy in the crystal lattice andbenefits the creep resistance property provided by the alloy. Cobaltalso contributes to the corrosion resistance provided by the alloy.Toward those ends, a preferred embodiment of the alloy contains at leastabout 0.5% cobalt. A second preferred embodiment contains at least about4% cobalt. A further embodiment contains at least about 9% cobalt. Toomuch cobalt can result in the precipitation of undesired phases such assigma phase (Co—Cr). Therefore, the alloy preferably contains not morethan about 10% cobalt. In one preferred embodiment the alloy containsnot more than about 6% cobalt. In another preferred embodiment the alloycontains not more than about 2% cobalt. In a further embodiment, thealloy contains not more than about 1% cobalt.

The elements carbon, silicon, zirconium, and boron may be present inthis alloy. However, those elements tend to depress the solidustemperature of the alloy and can segregate to the grain boundaries ofthe alloy matrix. In order to avoid or limit cracking on solidification,a narrow solidification temperature range is preferred. For thesereasons in the present alloy carbon is restricted to not more than about0.1%, silicon is restricted to not more than about 0.03%, zirconium islimited to not more than about 0.1%, and boron is limited to not morethan about 0.03%. In the alloy of this invention, the combined amount ofthe elements C, Si, Zr, and B is not more than about 0.15% andpreferably, not more than about 0.1%.

Up to about 1.5% iron may be present in the alloy of this invention.However, too much iron adversely affects the microstructural stabilityof the alloy and can combine with other elements to form undesiredsecondary phases such as sigma phase (Fe—Cr) and laves phase (e.g.,Fe₂Nb, Fe₂Ti, and Fe₂Mo). Therefore, the amount of iron is preferablyrestricted to not more than about 1.0%. Although it is not intentionallyadded, up to about 0.5% manganese may be present in this alloy as aresidual from alloying additions during melting.

The balance of the alloy is at least about 50% nickel. Preferably, thealloy contains about 50-75% nickel. The balance may also include minoramounts of inevitable impurity elements, such as phosphorus, sulfur,oxygen, and nitrogen, for example, not more than about 0.03% phosphorusand not more than about 0.01% each of sulfur oxygen and nitrogen.

The alloy according to this invention is preferably provided in powderform. Toward that end the alloy is vacuum induction melted and thenatomized to form fine powder. The resulting powder is then screened toappropriate size for the additive manufacturing process. Useful articlesare prepared using any known technique for metal powder consolidation,for example, selective laser powder bed fusion. After the article isfully formed, it is preferably processed using different combinationsand variations of hot isostatic pressing (HIP), solution heat treatment,quenching, and age hardening to develop the desired microstructure andproperties. Useful articles such as gas turbine components (e.g.,turbine blades) can also be made by direct casting of the alloy usingsuch casting techniques as directional solidification, equiaxed graincasting, and single crystal investment casting.

WORKING EXAMPLES

In order to demonstrate the desirable combination of properties providedby the alloy of this invention, a set of comparative examples weremelted, processed, and tested. Set forth in Table 1 below are the weightpercent compositions of four heats that were melted and tested.

TABLE 1 Al Co Cr Hf Mo Nb Ta Ti W Ni Example 1  4.05 5.01  9.79 0   2.96 3.86 7.93 0   0.86 Bal¹ Example 2  4.09 1.62  9.78 0    2.96 3.079.36 0   0.75 Bal¹ Example 3  3.95 0.11 5.8 1.18  0.99 0   9.37 0.527.37 Bal¹ CM247LC² 5.6 9.27 8.2 1.39 0.6 — 3.29 0.64 9.21 61.65 ¹Balanceincludes usual impurity elements. ²Also includes 0.03% Zr, 0.014% B, and0.082% C.

Examples 1, 2, and 3 were melted under vacuum and cast into ingots whichwere given a homogenization heat treatment after solidification.Examples 1 and 2 were solution heat treated at 2250° F. (1232° C.) forone hour and quenched in water. Example 3 was gas quenched from thehomogenization temperature in a chamber backfilled with 10 bar ofnitrogen. The CM247LC example was melted under vacuum, homogenized aftersolidification, solution treated at 2250° F. (1232° C.) for one hour andquenched in water. After quenching, specimens of each alloy were aged at1975° F. (1079° C.) for 4 hours followed by air cooling. The specimenswere then heat treated with a second aging at 1600° F. (871° C.) for 20hours and then air cooled.

The microstructures of Examples 1-3 after aging are shown in FIGS.1A-1C, respectively. The microstructure of an aged specimen of theCM247LC example is shown in FIG. 1D for comparison. The volume percentof γ′ in Examples 1-3 after aging is about 60 vol. % based on imageanalysis. The volume fraction of γ′ in Example 3 after aging is slightlyhigher than that of γ′ in Examples 1 and 2. The equivalent diameter ofthe γ′ precipitates in all 3 examples after ageing is about 300 to 500nm.

The lattice misfit parameters (6) of Examples 1-3 and CM247LC werecalculated using the THERMO-CALC software for a temperature of 1600° F.(871° C.). Actual lattice misfit parameters for Examples 1-3 and CM247LCwere measured by x-ray diffraction (XRD). The calculated and measuredvalues of 6 for each example are shown in Table 2 below.

TABLE 2 Example Calculated Lattice Misfit Measured Misfit 1 0.09 0.209 2−0.02 0.055 3 0.33 0.294 CM247LC −0.152 −0.035The results presented in Table 2 are graphed in FIG. 2 to show the goodcorrelation between the actual and calculated values.

Shown in FIGS. 3A-3C are optical images of specimens of Examples 1-3after being solution treated, aged, and then heated for 1000 h at 1600°F. (871° C.). It can be seen that a continuous external oxide layerforms on the alloy surface. All three of the examples aremicrostructurally stable without formation of detrimental phases duringexposure.

FIGS. 4A-4D show the melt pool structures in the alloy specimens after abead-on-plate test (the external surface is facing downward in thefigure). The sample of CM247LC (FIG. 4D) has cracks in the melt poolwhile the Examples 1-3 (FIGS. 4A, 4B, and 4C, respectively) are free ofcracks with all the power density and hatch-spacing combinations tested.

Additional heat treatment was carried out on the examples after thebead-on-plate test. The heat treatment consisted of heating the examplesat 1975° F. (1079° C.) for 4 hours, cooling in air, and then heating at1600° F. (871° C.) for 20 hours, followed by cooling in air. In thisprocedure, a laser is used to produce a weld-bead track on a metalplate. The results are shown in FIGS. 5A-5D. The aging heat treatmenteliminated the dendritic structure in the as-built condition and made itdifficult to discern the melt pool location. Compared to alloy CM247LC(FIG. 5D), which showed strain-age cracking, the three examples of thealloy of the present invention (FIGS. 5A, 5B, and 5C) show no sign ofcracks and which demonstrates that they are highly resistant tostrain-age cracking.

Based on the results from the bead-on-plate test, batches ofgas-atomized powder of the example alloys and the CM247LC alloy wereproduced. Additively manufactured samples were printed with a laserpowder bed fusion system. The processing parameters (including laserpower, scan speed, hatch spacing, etc.) were varied to provide aplurality of build samples. The as-built microstructures of the severalsamples are shown in FIGS. 6A to 6B. Each image was taken at 50×magnification. Thus, each image corresponds to a field of view about 2.6mm by 2.2 mm. In FIGS. 6A-6C, the laser power density increases fromleft to right while the hatch spacing decreases from top to bottom. FIG.6A shows the as-built microstructures of the CM247LC samples have severecracking issues for most test conditions. FIG. 6B, shows the as-builtmicrostructures of Example 1 and FIG. 6C shows the as-builtmicrostructures of Example 2. FIGS. 6B and 6C show good crack resistanceand demonstrate that the Examples 1 and 2 provide a large processingwindow in which to achieve good build quality with very low porosity andvery low crack density.

Samples with the lowest crack density and porosity were selected fromeach alloy and further processed with a post-build heat treatmentconsisting of solution and aging heat treatments. The samples weresolution treated at 2250° F. (1232° C.) for one hour, quenched in water,aged at 1975° F. (1079° C.) for 4 hours, air cooled, aged at 1600° F.(871° C.) for 20 hours, and then cooled in air. FIG. 7A shows asubstantial number of thin hairline cracks in the CM247LC material afteraging, while FIG. 7B and FIG. 7C show that the material of Example 1 andExample 2 remain crack free after post build heat treatment.

Additional examples were prepared by Argon gas atomization having theweight percent compositions shown in Table 3 below. The batches of alloypowder were consolidated and then processed to provide standardspecimens for tensile and stress rupture testing.

TABLE 3 Ex. Al Co Cr Mo Nb Ta W C B Zr Ni 4 4.0 5.2 9.9 3.0 4.0 7.9 1.20.077 0.015 0.013 Bal 5 4.1 1.6 9.7 3.0 3.1 8.9 3.8 0.038 0.007 0.010Bal

Tensile testing was performed on the specimens of Examples 4 and 5 atseveral temperatures in the range from room temperature up to 1800° F.(982° C.). The results of the tensile testing are set forth in Table 4below including the test temperature in degrees Celsius and the ultimatetensile strength and the yield strength in MPa. The results are alsoplotted in FIG. 8.

TABLE 4 Temp. Example 4 Example 5 (° C.) UTS(MPa) YS(MPa) UTS(MPa)YS(MPa) 21 1413.4 979.1 1461.7 1054.9 593 1285.9 865.3 1296.9 978.4 760946.7 903.9 1008.3 947.3 871 523.3 522.6 530.6 477.8 982 299.9 255.8228.9 —

Stress rupture testing was performed on the specimens of Examples 4 and5 at several temperatures in the range from 1300° F. to 1700° F. (704°C. to 927° C.). The results of the stress rupture testing are plotted inFIG. 9 based on the Larson-Miller parameter.

Cyclic oxidation testing was performed on specimens of Examples 4 and 5.The cyclic oxidation test was performed at 1800° F. (982° C.), with eachcycle made up of 0.25 hr ramping, a 1-hour hold at 1800° F. (982° C.),and 0.25 hr forced air cool. The results of the cyclic oxidation testingare set forth in Table 5 below including the number of cycles and thespecific weight change in g/cm².

TABLE 5 Specific Weight Change Cycles Example 4 Example 5 0 0 0 103.813(10⁻⁵) 3.472(10⁻⁴) 20 2.288(10⁻⁴⁾ 4.244(10⁻⁴) 50 7.626(10⁻⁵)7.716(10⁻⁴) 100 3.051(10⁻⁴) 8.488(10⁻⁴) 150 4.004(10⁻⁴) 1.138(10⁻³) 2005.148(10⁻⁴) 1.119(10⁻³) 300 4.004(10⁻⁴) 1.196(10⁻³) 400 2.097(10⁻⁴)1.389(10⁻³) 600 1.716(10⁻⁴) 1.447(10⁻³) 800 4.004(10⁻⁴) 1.678(10⁻³) 10003.813(10⁻⁵) 3.472(10⁻⁴)

FIG. 10 shows graphs of the data presented in Table 4. Both Example 4and Example 5 appear to be oxidation resistant since they did notexperience weight loss up to 1000 cycles.

The terms and expressions which are employed in this specification areused as terms of description and not of limitation. There is nointention in the use of such terms and expressions of excluding anyequivalents of the features shown and described or portions thereof. Itis recognized that various modifications are possible within theinvention described and claimed herein.

1. A high γ′ volume fraction Ni-base alloy optimized for laser andelectron beam additive manufacturing processes with improvedsolidification and strain-age cracking resistance and high strength andcreep resistance at elevated temperature, said alloy consistingessentially of, in weight percent, about C  0-0.1 Mn 0.5 max. Si  0-0.03 Cr  4-16 Fe 1.5 max. Mo 0-6 W 0-8 Co  0-15 Ti 0-2 Al 0.5-5.5 Nb0-6 Ta  7.5-14.5 Hf  0-2.0 Zr  0-0.1 Re 0-6 Ru 0-3 B   0-0.03

and the balance of the alloy is at least about 50% nickel and the usualimpurities found in commercial grades of Ni-base superalloys intendedfor the same use and service.
 2. The alloy claimed in claim 1 whichcontains, about Cr  9-11 Fe 1 max. Mo 2-5 W 0.5-2  Co 4-6 Ti  0-0.5 Al3-5 Nb 3-5 Ta 7.5-9.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 3. The alloy claimed in claim 1 which contains about Cr 9-11 Fe 1 max. Mo 2-5 W 0.5-2  Co 0.5-2  Ti  0-0.5 Al 3-5 Nb 2.5-4  Ta 8.5-10.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 4. The alloy claimed in claim 1 which contains about Cr 4-8Fe 1 max. Mo 0.5-1.5 W 6-8 Co 0-1 Ti 0-1 Al 3-5 Nb 0-1 Ta  8.5-10.5 Hf0.5-1.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 5. The alloy claimed in claim 1 which contains about Cr 9-11 Fe 1 max. Mo 2-5 W 3.5-4.5 Co 0.5-2  Ti  0-0.5 Al 3-5 Nb 2.5-4  Ta 8.5-10.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 6. The alloy claimed in claim 1 which contains about: Cr 8-10 Fe 1 max. Mo 2-5 W 4-6 Co  9-11 Ti  0-0.5 Al  4-5.5 Nb 0-1 Ta 9.5-11.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and the usualimpurities.
 7. The alloy claimed in claim 1 wherein Al+Ti is about 3 toabout 6 weight percent and W+Mo+Ru+Re is about 2 to about 12 weightpercent.
 8. The alloy claimed in claim 7 wherein C+B+Si+Zr is less thanabout 0.15 weight percent for reducing solidification range andimproving solidification cracking resistance.
 9. A high γ′ volumefraction Ni-base alloy optimized for laser and electron beam additivemanufacturing processes with improved solidification and strain-agecracking resistance and high strength and creep resistance at elevatedtemperature, said alloy consisting essentially of, in weight percent,about C  0-0.1 Mn 0.5 max.   Si   0-0.03 Cr  4-11 Fe 1 max. Mo 0-5 W0.5-8  Co  0-10 Ti 0-1 Al 3-5 Nb 0-5 Ta  7.5-11.5 Hf  0-1.5 Zr  0-0.1 Re0-5 Ru 0-3 B   0-0.03

and the balance of the alloy is at least about 50% nickel and the usualimpurities found in commercial grades of Ni-base superalloys intendedfor the same use and service.
 10. The alloy claimed in claim 9 whichcontains, about Cr  9-11 Fe 1 max. Mo 2-5 W 0.5-2  Co 4-6 Ti  0-0.5 Al3-5 Nb 3-5 Ta 7.5-9.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 11. The alloy claimed in claim 9 which contains about Cr 9-11 Fe 1 max. Mo 2-5 W 0.5-2  Co 0.5-2  Ti  0-0.5 Al 3-5 Nb 2.5-4  Ta 8.5-10.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 12. The alloy claimed in claim 9 which contains about Cr 4-8Fe 1 max. Mo 0.5-1.5 W 6-8 Co 0-1 Ti 0-1 Al 3-5 Nb 0-1 Ta  8.5-10.5 Hf0.5-1.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 13. The alloy claimed in claim 9 which contains about: Cr 9-11 Fe 1 max. Mo 2-5 W 3.5-4.5 Co 0.5-2  Ti  0-0.5 Al 3-5 Nb 2.5-4  Ta 8.5-10.5 Hf  0-0.5 Re 0-1 Ru 0-1

and the balance of the alloy is at least about 50% nickel and usualimpurities.
 14. The alloy claimed in claim 9 wherein Al+Ti is about 3.5to about 5.5 weight percent and W+Mo+Ru+Re is about 3 to about 10 weightpercent.
 15. The alloy claimed in claim 14 wherein C+B+Si+Zr is lessthan about 0.10 weight percent for reducing solidification range andimproving solidification cracking resistance.
 16. The alloy claimed inclaim 1 or claim 9 comprising a γ′ strengthening phase.
 17. The alloyclaimed in claim 16 wherein the γ′ phase has a volume fraction that isgreater than 35%.
 18. The alloy claimed in claim 17 wherein the γ′ phasehas a volume fraction that is greater than 50%.
 19. The alloy claimed inclaim 1 or claim 9 which has a γ′ solvus temperature 7 higher than 2050°F. (1121° C.).